JP5315888B2 - α-β type titanium alloy and method for melting the same - Google Patents
α-β type titanium alloy and method for melting the same Download PDFInfo
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本発明は、実用温度域において高強度で、疲労強度に優れると共に、高温度域における変形抵抗が小さく、熱間加工性に優れるチタン合金とその溶製方法に関するものである。 The present invention relates to a titanium alloy having high strength in a practical temperature range, excellent fatigue strength, low deformation resistance in a high temperature range, and excellent hot workability, and a method for melting the titanium alloy.
Ｔｉ−６Ａｌ−４Ｖ合金に代表されるα−β型チタン合金は、軽量かつ高強度で優れた耐食性を有していることから、航空機や自動車、スポーツ用品等のさまざまな分野で使用されている。しかしながら、高強度のα−β型チタン合金、特に、Ａｌを多く含有するα−β型チタン合金は、（α−β）２相温度域での熱間加工性に劣る、引張強度に比べて、疲労強度が低い等の欠点を有している。 Α-β type titanium alloys represented by Ti-6Al-4V alloy are lightweight and have high strength and excellent corrosion resistance, and are therefore used in various fields such as aircraft, automobiles and sports equipment. . However, high-strength α-β type titanium alloys, especially α-β type titanium alloys containing a large amount of Al, are inferior in hot workability in the (α-β) two-phase temperature range, compared with tensile strength. In addition, it has drawbacks such as low fatigue strength.
上記熱間加工性が劣るという欠点は、熱間での鍛造や圧延時の歩留りを低下させて、製造コストの上昇を招くため、チタン合金の用途拡大の大きな障害となっている。また、疲労強度が低いという欠点は、実用化されたチタン製部材の主な破損原因が疲労破壊であることを併せて考えれば、引張強度を単に高める材料開発だけでは、チタン合金の用途拡大に結びつかないことを示している。また、近年におけるＶ等レアメタルの価格の高騰によって、これらの元素を多量に含むチタン合金は、より高価なものとなってきている。 The disadvantage that the hot workability is inferior decreases the yield during hot forging and rolling, leading to an increase in production cost, and is a major obstacle to expanding the application of titanium alloys. In addition, the shortcoming of low fatigue strength is that the main cause of failure of titanium parts that have been put to practical use is fatigue failure. It shows that it is not connected. Further, in recent years, the price of rare metals such as V has soared, and titanium alloys containing a large amount of these elements have become more expensive.
上記問題に対しては、例えば、特許文献１には、Ｔｉ−６Ａｌ−４Ｖ合金の強度を高めると共に超塑性特性を向上させることによって熱間加工性を改善したＴｉ−４．５Ａｌ−３Ｖ−２Ｆｅ−２Ｍｏ合金が、また、特許文献２には、α安定化元素としてＡｌとＣ、β安定化元素としてＣｒとＦｅを適正量添加することにより、常温〜５００℃の実用温度域で高強度を有し、熱間加工性にも優れるＴｉ−４．５Ａｌ−４Ｃｒ−０．５Ｆｅ−０．２Ｃ合金が提案されている。
しかしながら、これらの合金はいずれも、ＦｅやＣｒといった凝固偏析を起こし易い元素を２〜４ｍａｓｓ％含有していることから、偏析を防止するために溶解速度を遅くする必要があり、製造コストが増大するという問題点がある。 However, these alloys all contain 2 to 4 mass% of elements such as Fe and Cr that are prone to solidification and segregation, so it is necessary to slow the dissolution rate to prevent segregation and increase the manufacturing cost. There is a problem of doing.
そこで本発明は、上記事情に鑑みてなされたものであり、その目的は、室温から４００℃の実用温度域において高強度α−β型チタン合金として広く用いられているＴｉ−６Ａｌ−４Ｖ合金に匹敵する強度を有するとともに疲労強度にも優れ、さらに８００℃前後の高温における熱間加工性にも優れるチタン合金を低コストで提供することにある。 Therefore, the present invention has been made in view of the above circumstances, and the object of the present invention is to a Ti-6Al-4V alloy widely used as a high-strength α-β type titanium alloy in a practical temperature range from room temperature to 400 ° C. The object is to provide a titanium alloy having comparable strength, excellent fatigue strength, and excellent hot workability at a high temperature of about 800 ° C. at a low cost.
発明者らは、上記課題を解決するべく、チタン合金に添加される各種合金元素が上記特性に及ぼす影響について鋭意検討を重ねた。その結果、Ｔｉ−６Ａｌ−４Ｖ合金に匹敵する強度と、優れた疲労強度を兼備するためには、後述するアルミ当量Ａｌｅｑを適正範囲に制御する必要があること、また、優れた熱間加工性を確保するためには、後述するβ変態点温度Ｔβを適正範囲に制御する必要があること、さらに、疲労強度を改善するには、合金溶製時のＮ添加原料を適正化する必要があることを見出し、本発明を完成させた。 In order to solve the above-mentioned problems, the inventors have conducted intensive studies on the influence of various alloy elements added to the titanium alloy on the above characteristics. As a result, in order to combine strength comparable to that of Ti-6Al-4V alloy and excellent fatigue strength, it is necessary to control the aluminum equivalent Al eq described later within an appropriate range, and excellent hot working In order to secure the properties, it is necessary to control the β transformation point temperature T β described later within an appropriate range, and in order to improve fatigue strength, it is necessary to optimize the N-added raw material at the time of alloy melting As a result, the present invention has been completed.
すなわち本発明は、Ａｌ：３．０〜５．０ｍａｓｓ％、Ｖ：１．０〜３．０ｍａｓｓ％、Ｆｅ：１．０〜１．８ｍａｓｓ％、Ｍｏ：０．９〜１．７ｍａｓｓ％、Ｏ：０．０５〜０．２５ｍａｓｓ％およびＮ：０．０２〜０．１５ｍａｓｓ％を含有し、残部がＴｉおよび不可避的不純物からなり、下記（１）式；
Ａｌｅｑ（ｍａｓｓ％）＝［Ａｌ］＋１０×［Ｏ］＋２７．７×［Ｎ］ ・・・（１）
で表されるアルミ当量Ａｌｅｑが４〜８ｍａｓｓ％、下記（２）式；
Ｔβ（℃）＝８８６＋１４７．７×［Ｏ］＋２９４．３×［Ｎ］＋２０．４×［Ａｌ］−１９．８×［Ｆｅ］−１３．１×［Ｖ］−１０．３×［Ｍｏ］ ・・・（２）
で表されるβ変態温度Ｔβが８８０〜９８０℃であるα−β型チタン合金である。
That is, the present invention is Al: 3.0-5.0 mass%, V: 1.0-3.0 mass%, Fe: 1.0-1.8 mass%, Mo: 0.9-1.7 mass%, O : 0.05 to 0.25 mass % and N: 0.02 to 0.15 mass% , the balance is composed of Ti and inevitable impurities, the following formula (1);
Al eq (mass%) = [Al] + 10 × [O] + 27.7 × [N] (1)
The aluminum equivalent Al eq represented by the formula is 4 to 8 mass%, the following formula (2);
T β (° C.) = 886 + 147.7 × [O] + 294.3 × [N] + 20.4 × [Al] −19.8 × [Fe] −13.1 × [V] −10.3 × [Mo ] (2)
Is an α-β type titanium alloy having a β transformation temperature T β of 880 to 980 ° C.
また、本発明は、上記のＴｉ合金を溶製するに当たり、溶解原料にＦｅ−Ｖ−Ｎ母合金を使用してＮを添加することを特徴とするα−β型チタン合金の溶製方法を提案する。 In addition, the present invention provides a method for melting an α-β type titanium alloy characterized by adding N using a Fe-V-N master alloy as a melting raw material when melting the Ti alloy. suggest.
本発明によれば、室温から４００℃の実用温度域における高い強度と、優れた疲労強度を有するとともに、８００℃以上の熱間加工性にも優れるα−β型チタン合金を低コストで提供することができる。したがって、本発明によれば、製造コストの低減やチタン合金への信頼性の向上を通じて、航空機部品や自動車部品、スポーツ用品の分野のみならず、それ以外の分野にも用途拡大が可能となるので、産業上その効果は大である。 According to the present invention, an α-β type titanium alloy having high strength in a practical temperature range from room temperature to 400 ° C. and excellent fatigue strength and excellent hot workability at 800 ° C. or higher is provided at low cost. be able to. Therefore, according to the present invention, it is possible to expand the application not only to the field of aircraft parts, automobile parts, and sporting goods, but also to other fields through reduction in manufacturing cost and improvement in reliability of titanium alloys. The effect on the industry is great.
本発明のチタン合金における成分組成の限定理由について説明する。
Ａｌ：３．０〜５．０ｍａｓｓ％
チタン合金の熱間加工は、熱間鍛造か熱間圧延あるいはその両者を組み合わせて行われるのが普通である。しかし、熱間加工中に温度が８００℃以下に低下すると、変形抵抗が増大するとともに、割れ等の欠陥が発生し易くなり、製造性が著しく低下する。この製造性は、Ａｌの含有量と密接な関係がある。即ち、Ａｌは、（α−β）２相組織を得るためのα相安定化元素として添加される元素であり、強度の上昇にも寄与する。しかし、Ａｌ含有量が３．０ｍａｓｓ％未満では、十分な強度が得られない。一方、Ａｌ含有量が５．０ｍａｓｓ％を超えると、熱間変形抵抗が増大すると共に、割れ感受性が著しく増大して製造性が悪化し、製造コストの増大に繋がる。よって、Ａｌの含有量は３．０〜５．０ｍａｓｓ％の範囲とする。
The reason for limiting the component composition in the titanium alloy of the present invention will be described.
Al: 3.0-5.0 mass%
The hot working of a titanium alloy is usually performed by hot forging, hot rolling, or a combination of both. However, when the temperature is lowered to 800 ° C. or lower during hot working, deformation resistance is increased and defects such as cracks are likely to occur, resulting in a significant reduction in manufacturability. This manufacturability is closely related to the Al content. That is, Al is an element added as an α-phase stabilizing element for obtaining an (α-β) two-phase structure, and contributes to an increase in strength. However, if the Al content is less than 3.0 mass%, sufficient strength cannot be obtained. On the other hand, when the Al content exceeds 5.0 mass%, hot deformation resistance increases, crack sensitivity is remarkably increased, manufacturability is deteriorated, and manufacturing cost is increased. Therefore, the Al content is in the range of 3.0 to 5.0 mass%.
Ｖ：１．０〜３．０ｍａｓｓ％
Ｖは、（α−β）２相組織を得るためのβ相安定化元素として添加される元素であり、Ｔｉとの間で脆化相である金属間化合物を形成することなく主にβ相に固溶して、強度の上昇に寄与する。しかし、Ｖは、高融点金属であると共に、近年、特に価格が高騰しているレアメタルの一つであることから、その含有量を抑えることが好ましい。一方、現在、最も広く使用されているＴｉ-６Ａｌ-４Ｖ合金は、そのスクラップの流通量も多く、スクラップのリサイクルが可能となることから、Ｖをある程度含有させても溶解コストの大幅な上昇とはならない。よって、本発明では、Ｖは１．０〜３．０ｍａｓｓ％の範囲で含有させる。
V: 1.0-3.0 mass%
V is an element added as a β-phase stabilizing element for obtaining an (α-β) two-phase structure, and is mainly a β-phase without forming an intermetallic compound that is an embrittled phase with Ti. It contributes to the increase in strength. However, since V is a refractory metal and is one of rare metals whose prices have been rising in recent years, it is preferable to suppress the content thereof. On the other hand, currently the most widely used Ti-6Al-4V alloy has a large amount of scrap and can be recycled. Therefore, even if it contains V to some extent, the melting cost will increase significantly. Must not. Therefore, in the present invention, V is contained in the range of 1.0 to 3.0 mass%.
Ｆｅ：１．０〜１．８ｍａｓｓ％
Ｆｅは、β安定化元素であり、主にβ相に固溶し、合金強度の上昇に寄与する。また、α相に比べて熱間加工性のよいβ相の体積分率を増大させる作用があることと、β相中の拡散速度が大きいことが相俟って、チタン合金の熱間変形抵抗を減少させるのに寄与する。しかし、含有量が１．０ｍａｓｓ％未満では、その効果が十分ではなく、一方、含有量が過大になると、Ｔｉとの間で脆化相である金属間化合物（ＴｉＦｅ）を形成し、さらに、溶解・凝固時にβフレックと呼ばれる偏析相を生成して、合金の機械的性質、特に延性、靭性の低下をもたらす。よって、本発明では、Ｆｅは１．０〜１．８ｍａｓｓ％の範囲とする。
Fe: 1.0-1.8 mass%
Fe is a β-stabilizing element, and is mainly dissolved in the β-phase and contributes to an increase in alloy strength. Also, it has the effect of increasing the volume fraction of the β phase, which has better hot workability than the α phase, and the high diffusion rate in the β phase. Contributes to reducing However, if the content is less than 1.0 mass%, the effect is not sufficient. On the other hand, if the content is excessive, an intermetallic compound (TiFe) that is an embrittlement phase with Ti is formed. When melted and solidified, a segregation phase called β-flex is generated, and the mechanical properties, particularly ductility and toughness of the alloy are lowered. Therefore, in this invention, Fe is taken as the range of 1.0-1.8 mass%.
Ｍｏ：０．９〜１．７ｍａｓｓ％
Ｍｏは、Ｖと同様、β安定化元素であり、主にβ相に固溶して強度の上昇に寄与する。また、α相に比べ熱間加工性のよいβ相の体積分率を増大させることにより、鍛造、圧延等の塑性加工性を向上する効果がある。しかし、Ｍｏの含有量が０．９ｍａｓｓ％未満では、その効果が十分でなく、一方、１．７ｍａｓｓ％を超えると、Ｍｏが重い元素であることから合金の比重を大きくし、高比強度であるというチタン合金の特長を損なう。さらに、Ｍｏは、チタン中での拡散速度が小さいために、熱間加工時の変形抵抗を増大させる。よって、Ｍｏは０．９〜１．７ｍａｓｓ％の範囲で添加する。
Mo: 0.9 to 1.7 mass%
Mo, like V, is a β-stabilizing element and mainly contributes to an increase in strength by solid solution in the β phase. Further, by increasing the volume fraction of the β phase, which has better hot workability than the α phase, there is an effect of improving plastic workability such as forging and rolling. However, if the Mo content is less than 0.9 mass%, the effect is not sufficient. On the other hand, if it exceeds 1.7 mass%, Mo is a heavy element, so the specific gravity of the alloy is increased and the specific strength is high. The feature of titanium alloy is impaired. Furthermore, since Mo has a low diffusion rate in titanium, it increases deformation resistance during hot working. Therefore, Mo is added in the range of 0.9 to 1.7 mass%.
Ｏ：０．０５〜０．２５ｍａｓｓ％
Ｏは、α相に固溶して強度を上昇させる作用を有する。しかしながら、Ｏの含有量が０．０５ｍａｓｓ％未満では、強度上昇への寄与が十分でなく、所望の強度が得られない。一方、Ｏが０．２５ｍａｓｓ％を超えると、室温での延性や靭性が低下し、また加工性も低下するので好ましくない。よって、Ｏは０．０５〜０．２５ｍａｓｓ％の範囲とする。
O: 0.05-0.25 mass%
O has a function of increasing the strength by dissolving in the α phase. However, if the O content is less than 0.05 mass%, the contribution to the strength increase is not sufficient, and the desired strength cannot be obtained. On the other hand, if O exceeds 0.25 mass%, ductility and toughness at room temperature are lowered, and workability is also lowered, which is not preferable. Therefore, O is in the range of 0.05 to 0.25 mass%.
Ｎ：０．０２〜０．１５ｍａｓｓ％
Ｎは、Ｏと同様、α相に固溶して強度を上昇させる作用を有する元素である。通常、Ｎは、不可避的に０．００１〜００１ｍａｓｓ％程度混入するが、上記効果を得るためには、Ｎを積極的に添加するのが好ましい。この目的で添加する場合、Ｎの添加量が０．０２ｍａｓｓ％未満では、強度上昇への寄与が十分でなく、所望の強度が得られない。一方、Ｎ含有量が０．１５ｍａｓｓ％を超えると、室温での延性や靭性が低下する。さらに、チタン合金を溶製する際、Ｎの含有量が多くなりすぎると、Ｎ添加原料の融け残りが問題になることがある。よって、Ｎを添加する場合は０．０２〜０．１５ｍａｓｓ％の範囲とするのが好ましい。
N: 0.02-0.15 mass%
N, like O, is an element that has the effect of increasing the strength by dissolving in the α phase. Normally, N is inevitably mixed in an amount of about 0.001 to 001 mass%. However, in order to obtain the above effect, it is preferable to add N positively. When added for this purpose, if the addition amount of N is less than 0.02 mass%, the contribution to strength increase is not sufficient, and the desired strength cannot be obtained. On the other hand, if the N content exceeds 0.15 mass%, the ductility and toughness at room temperature decrease. Further, when the titanium alloy is melted, if the N content becomes too large, the unmelted N-added raw material may become a problem. Therefore, when adding N, it is preferable to set it as the range of 0.02-0.15 mass%.
本発明のチタン合金は、上記成分組成の規制に加えてさらに、以下に説明するアルミ当量Ａｌｅｑおよびβ変態点温度Ｔβが、本発明の条件を満たすことが必要である。
アルミ当量Ａｌｅｑ：４．０〜８．０ｍａｓｓ％
アルミ当量Ａｌｅｑは、チタン合金に含まれるα相安定化元素の量を表すパラメータの一つであり、α相を強化する程度を示している。このＡｌｅｑは、下記（１）式；
Ａｌｅｑ（ｍａｓｓ％）＝［Ａｌ］＋１０×［Ｏ］＋２７．７×［Ｎ］ ・・・（１）
ここで、［Ａｌ］、［Ｏ］および［Ｎ］は、それぞれの元素の含有量（ｍａｓｓ％）
で定義される。Ａｌｅｑの値が、４ｍａｓｓ％未満では、α相の強化が十分でなく、所望の強度が得られない。一方、Ａｌｅｑが８ｍａｓｓ％を超えると、熱間加工性を低下させるだけでなく、機械的性質、特に疲労特性を著しく低下させる。よって、Ａｌｅｑは、４．０〜８．０ｍａｓｓ％の範囲に制御する必要がある。
The titanium alloy of the present invention is required to satisfy the conditions of the present invention in addition to the regulation of the component composition described above, and further, an aluminum equivalent Al eq and a β transformation point temperature T β described below.
Aluminum equivalent Al eq : 4.0 to 8.0 mass%
Aluminum equivalent Al eq is one of the parameters representing the amount of the α-phase stabilizing element contained in the titanium alloy, and indicates the degree to which the α-phase is strengthened. This Al eq is the following formula (1):
Al eq (mass%) = [Al] + 10 × [O] + 27.7 × [N] (1)
Here, [Al], [O] and [N] are the content of each element (mass%)
Defined by If the value of Al eq is less than 4 mass%, the α phase is not sufficiently strengthened and the desired strength cannot be obtained. On the other hand, when Al eq exceeds 8 mass%, not only hot workability is lowered but also mechanical properties, particularly fatigue properties, are remarkably lowered. Therefore, Al eq needs to be controlled in the range of 4.0 to 8.0 mass%.
β変態点温度Ｔβ：８８０〜９８０℃
β変態点温度Ｔβは、高温度側のβ単相からα相が平衡的に生成を開始する温度であり、合金を構成する各元素と下記（２）式；
Ｔβ（℃）＝８８６＋１４７．７×［Ｏ］＋２９４．３×［Ｎ］＋２０．４×［Ａｌ］−１９．８×［Ｆｅ］−１３．１×［Ｖ］−１０．３×［Ｍｏ］ ・・・（２）
ここで、［Ｏ］、［Ｎ］、［Ａｌ］、［Ｆｅ］、［Ｖ］および［Ｍｏ］は、それぞれの元素の含有量（ｍａｓｓ％）
の関係がある。このＴβは、チタン合金におけるβ相の安定度を示すパラメータとしても用いられ、チタン合金を熱間加工する際には、Ｔβ−１００℃前後の温度に加熱するのが通常である。Ｔβが９８０℃超であると、熱間加工のための加熱温度が高くなり過ぎて、加熱時に材料表面が過度に酸化したり、結晶粒が粗大化したりして、熱間加工時に割れが発生し易くなる。一方、この温度が８８０℃未満であると、熱間加工の際の加熱温度が低くなり過ぎて、やはり割れが発生し易くなると共に、β相が安定化し過ぎて、溶体化時効処理で時効硬化させるのに要する時間が長くなり過ぎる。よって、β変態点温度Ｔβは、８８０〜９８０℃の範囲となるよう合金成分を調整する必要がある。
β transformation temperature T β : 880-980 ° C
β transformation point temperature T β is a temperature at which α phase starts to form in equilibrium from β single phase on the higher temperature side, each element constituting the alloy and the following formula (2):
T β (° C.) = 886 + 147.7 × [O] + 294.3 × [N] + 20.4 × [Al] −19.8 × [Fe] −13.1 × [V] −10.3 × [Mo ] (2)
Here, [O], [N], [Al], [Fe], [V] and [Mo] are the content of each element (mass%).
There is a relationship. This T beta, also used as a parameter indicating the stability of the beta phase in the titanium alloy, when the titanium alloy to hot working, it is usual to heat the T beta -100 ° C. temperature of about. If Tβ exceeds 980 ° C., the heating temperature for hot working becomes too high, the surface of the material is excessively oxidized during heating, the crystal grains become coarse, and cracks occur during hot working. It tends to occur. On the other hand, if this temperature is less than 880 ° C., the heating temperature during hot working becomes too low, and cracking is likely to occur, and the β phase is too stabilized, so that age hardening is achieved by solution aging treatment. It takes too long to make it happen. Therefore, it is necessary to adjust the alloy components so that the β transformation point temperature T β is in the range of 880 to 980 ° C.
本発明のチタン合金は、上記成分以外の残部は、Ｔｉおよび不可避的不純物である。なお、本発明のチタン合金は、本発明の作用効果を害さない範囲であれば、上記成分以外の成分の含有を拒むものではなく、例えば、Ｃ：０．０８ｍａｓｓ％以下、Ｈ：０．０５ｍａｓｓ％以下、より好ましくは０．０１５ｍａｓｓ％以下を含有することができる。 In the titanium alloy of the present invention, the balance other than the above components is Ti and inevitable impurities. In addition, the titanium alloy of the present invention does not refuse to contain components other than the above components as long as the effects of the present invention are not impaired. For example, C: 0.08 mass% or less, H: 0.05 mass % Or less, more preferably 0.015 mass% or less.
上記成分組成を満たす本発明のチタン合金は、従来のＴｉ−６Ａｌ−４Ｖ合金と同等の強度を有し、かつこれよりも高い延性と優れた疲労強度を有する。さらに、従来のＴｉ−６Ａｌ−４Ｖ合金に比べて、優れた熱間加工性を有していることから、従来合金と同じ工程、同じ設備で製造することが可能である。 The titanium alloy of the present invention satisfying the above component composition has the same strength as a conventional Ti-6Al-4V alloy, and has higher ductility and superior fatigue strength. Furthermore, since it has excellent hot workability as compared with the conventional Ti-6Al-4V alloy, it can be manufactured in the same process and the same equipment as the conventional alloy.
次に、Ｎを添加したチタン合金を溶製する場合の本発明の溶製方法について説明する。
本発明の溶製方法において、チタン合金にＮを添加する場合は、溶解原料としてＦｅ−Ｖ−Ｎ合金を用いて行うことを特徴とする。
一般に、溶解したチタンあるいはチタン合金に酸素を添加する場合には、二酸化チタン（ＴｉＯ２）の粉末が用いられている。それは、ＴｉＯ２の融点が約１６４０℃と、チタンの融点である約１６６８℃に近く、溶解が比較的容易であるからである。しかし、従来から窒素の添加に用いられているチタンと窒素の化合物である窒化チタン（ＴｉＮ）は、融点が約３２９０℃と非常に高い。そのため、ＴｉＮを用いてＮ添加を行った場合、たとえ超微細紛のＴｉＮを用いたとしても、通常の溶解方法では融け残りが発生するおそれがある。そこで、本発明は、低融点のＦｅ−Ｖ−Ｎ合金、具体的には、４３ｍａｓｓ％Ｆｅ-５０ｍａｓｓ％Ｖ-７ｍａｓｓ％Ｎ（融点：約１５２０℃）をＮ源として用いて、Ｎの添加を行う。これにより、従来と同じ真空アーク溶解法（ＶＡＲ）を用いた溶解方法でも、Ｎ偏析のないチタン合金を得ることができる。
Next, the melting method of the present invention when melting a titanium alloy added with N will be described.
In the melting method of the present invention, when adding N to a titanium alloy, it is performed using an Fe-V-N alloy as a melting raw material.
Generally, when oxygen is added to dissolved titanium or a titanium alloy, titanium dioxide (TiO 2 ) powder is used. This is because the melting point of TiO 2 is about 1640 ° C., which is close to about 1668 ° C., which is the melting point of titanium, and is relatively easy to dissolve. However, titanium nitride (TiN), which is a compound of titanium and nitrogen conventionally used for the addition of nitrogen, has a very high melting point of about 3290 ° C. For this reason, when N is added using TiN, there is a possibility that unmelted residue may be generated by a normal melting method even if TiN as an ultrafine powder is used. Therefore, the present invention uses a low melting point Fe—V—N alloy, specifically, 43 mass% Fe-50 mass% V-7 mass% N (melting point: about 1520 ° C.) as an N source, and the addition of N Do. As a result, a titanium alloy having no N segregation can be obtained even by a melting method using the same vacuum arc melting method (VAR) as in the prior art.
表１に示した成分組成を有する本発明に適合する発明合金（Ｎｏ．Ａ１〜Ａ９（ただし、Ｎｏ．Ａ１〜Ａ５は参考合金））、従来から公知の従来合金（Ｎｏ．Ｂ１〜Ｂ３）、および比較合金（Ｎｏ．Ｃ１〜Ｃ９）を、真空アーク溶解炉（ＶＡＲ）を用いて溶製し、鋳造して鋳塊とし、この鋳塊を熱間鍛造し、熱間圧延して直径が１８ｍｍφの丸棒に仕上げた。なお、上記熱間鍛造する際の加熱温度は、一次鍛造はそれぞれの合金が有するβ変態点温度以上である１１００℃、二次鍛造はそれぞれの合金が有するβ変態点温度より７０℃低い温度に設定した。また、丸棒に熱間圧延する際の加熱温度は、それぞれの合金が有するβ変態点温度より１２０℃低い温度とし、圧下量は断面減少率で約８０〜９５％とした。その後、上記丸棒に７２０℃で１ｈｒの焼鈍を施したのち、下記に試験に供した。 Invention alloys (No. A1 to A9 (where No. A1 to A5 are reference alloys) ) suitable for the present invention having the composition shown in Table 1, conventionally known alloys (No. B1 to B3), And the comparative alloys (No. C1 to C9) were melted using a vacuum arc melting furnace (VAR), cast into an ingot, the ingot was hot forged, hot rolled, and a diameter of 18 mmφ Finished in a round bar. In addition, the heating temperature at the time of hot forging is 1100 ° C. which is higher than the β transformation point temperature of each alloy in primary forging, and 70 ° C. lower than the β transformation point temperature of each alloy in secondary forging. Set. Moreover, the heating temperature at the time of hot rolling to a round bar was set to 120 ° C. lower than the β transformation point temperature of each alloy, and the reduction amount was about 80 to 95% in terms of the cross-sectional reduction rate. Thereafter, the round bar was annealed at 720 ° C. for 1 hr, and then subjected to the following test.
＜引張試験＞
上記のようにして得たそれぞれの丸棒から、丸棒の長手方向が試験片の長さ方向に平行になるように、ＡＳＴＭ Ｅ８に規定された平行部６．２５ｍｍφ、標点間距離２５ｍｍの引張試験片を採取し、ＡＳＴＭ Ｅ８に準拠して室温で引張試験を行い、０．２％耐力（０．２％ＰＳ）、引張強さ（ＵＴＳ）、伸び（Ｅｌ）および絞り（ＲＡ）を測定した。
＜熱間変形抵抗＞
丸棒の長手方向が試験片の長さ方向に平行になるようにして直径８ｍｍφ×長さ１２ｍｍの円柱状試験片を採取し、高温圧縮試験機（富士電波工機社製「サーメックマスターＺ」）を用いて、真空中で、高周波で８００℃に加熱後、歪速度１０（１／秒）で単軸圧縮加工する熱間圧縮鍛造模擬試験を行い、長さ方向に５０％圧縮した時の真応力を測定し、熱間変形抵抗を評価した。
＜疲労試験＞
上記引張試験および熱間変形抵抗を測定したうちのいくつかの丸棒から、ＪＩＳ Ｚ２２７４「金属材料の回転曲げ疲れ試験方法」に規定された１号試験片を採取し、同じくＪＩＳ Ｚ２２７４に準拠して回転曲げ疲れ試験に供してＳ−Ｎ線図を求め、繰り返し数が１０７サイクルでも破断しない繰り返し応力で定義する疲労限度を求めた。
<Tensile test>
From each of the round bars obtained as described above, a parallel part of 6.25 mmφ defined by ASTM E8 and a distance between gauge points of 25 mm are set so that the longitudinal direction of the round bar is parallel to the length direction of the test piece. Tensile test specimens are collected and subjected to a tensile test at room temperature in accordance with ASTM E8, and 0.2% proof stress (0.2% PS), tensile strength (UTS), elongation (El) and drawing (RA) are determined. It was measured.
<Hot deformation resistance>
A cylindrical test piece having a diameter of 8 mmφ × length of 12 mm was taken so that the longitudinal direction of the round bar was parallel to the length direction of the test piece, and a high-temperature compression tester (“Cermec Master Z” manufactured by Fuji Electric Koki Co., Ltd.) was collected. )), A hot compression forging simulation test was performed in which a single-axis compression process was performed at a strain rate of 10 (1 / second) after heating to 800 ° C. at high frequency in a vacuum, and 50% compression was performed in the length direction. The true stress was measured and the hot deformation resistance was evaluated.
<Fatigue test>
Sample No. 1 specified in JIS Z2274 “Rotating Bending Fatigue Test Method for Metallic Materials” was collected from some of the round bars from which the tensile test and hot deformation resistance were measured, and in accordance with JIS Z2274. was subjected to a bending fatigue test rotation seek S-N line shown here with, to determine the fatigue limit as defined repetitive stress repetition number is not broken even at 10 7 cycles.
上記試験の結果をまとめて表２に示した。また、表２には、熱間変形抵抗を疲労限度で除した値（表中のＡ／Ｂ）も併記した。これらの結果から、本発明のチタン合金（Ｎｏ．Ａ１〜Ａ９（ただし、Ｎｏ．Ａ１〜Ａ５は参考合金））は、従来のＴｉ−６Ａｌ−４Ｖ合金（Ｎｏ．Ｂ１）とほぼ同等の引張強さを有すること、また、疲労限度は、いずれも従来合金を上回っていることがわかる。また、熱間変形抵抗は、従来合金よりも小さく、熱間変形抵抗を疲労限度で除した値（Ａ／Ｂ）も、従来合金や比較合金では０．４９以上であるのに対して、本発明合金は、０．３０〜０．３４と小さく、熱間加工性にも優れていることがわかる。また、図１は、上記疲労試験に供した材料の、疲労限度と熱間変形抵抗との関係を示したものである。これらの結果から、本発明のチタン合金は、疲労限度が高い割に、熱間加工がしやすい材料であることがわかる。 The results of the above tests are summarized in Table 2. Table 2 also shows the value obtained by dividing the hot deformation resistance by the fatigue limit (A / B in the table). From these results, the titanium alloys of the present invention (No. A1 to A9 (where No. A1 to A5 are reference alloys) ) have substantially the same tensile strength as the conventional Ti-6Al-4V alloy (No. B1). It can be seen that the fatigue limit is higher than that of conventional alloys. Moreover, the hot deformation resistance is smaller than that of the conventional alloy, and the value (A / B) obtained by dividing the hot deformation resistance by the fatigue limit is 0.49 or more in the conventional alloy and the comparative alloy, whereas It can be seen that the invention alloy is as small as 0.30 to 0.34 and is excellent in hot workability. FIG. 1 shows the relationship between the fatigue limit and hot deformation resistance of the material subjected to the fatigue test. From these results, it can be seen that the titanium alloy of the present invention is a material that is easily hot-worked despite its high fatigue limit.
表１のＮｏ．２２，２３（Ｄ１，Ｄ２）に示した、ほぼ同じ成分組成を有するチタン合金を、真空アーク溶解法を用いて溶製し、鋳造し、直径２００ｍｍφの鋳塊とした。なお、両合金の溶製条件は、Ｎ添加原料としては、Ｎｏ．Ｄ１の合金はＦｅ−Ｖ−Ｎ合金を、また、Ｎｏ．Ｄ２の合金はＴｉＮ粉末を用いたこと以外は同一とした。次いで、上記鋳塊を１１００℃に加熱し、熱間鍛造して直径１２０ｍｍφの一次ビレットとし、その後、８７０℃に加熱し、鍛造して直径８０ｍｍφの二次ビレットとした。これをさらに８２０℃に加熱し、熱間圧延して直径１８ｍｍφの丸棒としたのち、７２０℃で１ｈｒの焼鈍を施した。
次いで、上記丸棒のそれぞれから、平行部が８ｍｍφ×２０ｍｍのダンベル型の疲労試験片を各１０本ずつ採取し、油圧サーボ式疲労試験機を用いて最大応力６１０Ｎ／ｍｍ２、応力比（最小応力に対する最大応力の比）−１の条件で両振り応力を加える疲労試験に供し、破断するまでのサイクル数（疲労寿命）を測定した。
No. in Table 1 Titanium alloys having substantially the same composition as shown in Nos. 22 and 23 (D1, D2) were melted and cast using a vacuum arc melting method to form an ingot having a diameter of 200 mmφ. The melting conditions of both alloys are No. as the N-added raw material. The alloy of D1 is an Fe-V-N alloy, and The alloy of D2 was the same except that TiN powder was used. Next, the ingot was heated to 1100 ° C. and hot forged to a primary billet having a diameter of 120 mmφ, then heated to 870 ° C. and forged to a secondary billet having a diameter of 80 mmφ. This was further heated to 820 ° C., hot-rolled into a round bar having a diameter of 18 mmφ, and then annealed at 720 ° C. for 1 hr.
Next, 10 dumbbell-shaped fatigue test pieces each having a parallel portion of 8 mmφ × 20 mm were sampled from each of the round bars, and a maximum stress of 610 N / mm 2 and a stress ratio (minimum) were obtained using a hydraulic servo fatigue tester. The ratio of maximum stress to stress) -1 was subjected to a fatigue test in which a swing stress was applied, and the number of cycles (fatigue life) until rupture was measured.
上記試験の結果を表３に示した。本発明に従いＦｅ−Ｖ−Ｎ合金を用いてＮ添加したＮｏ．Ｄ１の合金では、いずれの試験片でもほぼ一定の安定した疲労寿命が得られている。これに対して、ＴｉＮを用いてＮ添加した材料Ｎｏ．Ｄ２は、一部の試験片に疲労寿命が極端に短いものが見られ、平均疲労寿命がＤ１の約３／４で、標準偏差が約３倍あり、信頼性に劣る結果となった。そこで、Ｄ１の試験片の破面を肉眼で観察したところ、疲労寿命が短かった試験片は、いずれも疲労き裂の起点が破断面内であるのに対し、疲労寿命が長かった試験片は、いずれも試験片表面から疲労き裂が発生していた。そこで、疲労寿命が短かった試験片の破断面をさらに走査型電子顕微鏡で詳細に観察したところ、疲労き裂の起点と思われる箇所に、窒素が濃化した析出物が認められた。この析出物を分析したところ、Ｎ添加に用いた未溶解の窒化物（ＴｉＮ）であることが確認された。 The results of the above test are shown in Table 3. In accordance with the present invention, N-added No. No. was added using an Fe—V—N alloy. In the alloy of D1, almost constant and stable fatigue life is obtained in any specimen. On the other hand, the material No. in which N is added using TiN. In D2, some of the test pieces had extremely short fatigue life, the average fatigue life was about 3/4 of D1, and the standard deviation was about three times, resulting in poor reliability. Therefore, when the fracture surface of the D1 test piece was observed with the naked eye, all of the test pieces with a short fatigue life had a fatigue crack starting point within the fracture surface, whereas the test pieces with a long fatigue life were In both cases, fatigue cracks occurred from the surface of the test piece. Then, when the fracture surface of the test piece with a short fatigue life was further observed with a scanning electron microscope, precipitates enriched with nitrogen were observed at the places considered to be fatigue crack initiation points. When this precipitate was analyzed, it was confirmed to be undissolved nitride (TiN) used for N addition.
本発明のチタン合金は、航空機や自動車、スポーツ用品等に加えてさらに、高速鉄道や高速船舶、ロケット、医療機器、蒸気タービンブレードなどの分野にも好適に用いることができる。 The titanium alloy of the present invention can be suitably used in fields such as high-speed railways, high-speed ships, rockets, medical equipment, and steam turbine blades in addition to aircraft, automobiles, and sporting goods.
Claims (2)
Ｖ：１．０〜３．０ｍａｓｓ％、
Ｆｅ：１．０〜１．８ｍａｓｓ％、
Ｍｏ：０．９〜１．７ｍａｓｓ％、
Ｏ：０．０５〜０．２５ｍａｓｓ％および
Ｎ：０．０２〜０．１５ｍａｓｓ％を含有し、
残部がＴｉおよび不可避的不純物からなり、下記（１）式で表されるアルミ当量Ａｌｅｑが４〜８ｍａｓｓ％、下記（２）式で表されるβ変態温度Ｔβが８８０〜９８０℃であるα−β型チタン合金。
記
Ａｌｅｑ（ｍａｓｓ％）＝［Ａｌ］＋１０×［Ｏ］＋２７．７×［Ｎ］ ・・・（１）
Ｔβ（℃）＝８８６＋１４７．７×［Ｏ］＋２９４．３×［Ｎ］＋２０．４×［Ａｌ］−１９．８×［Ｆｅ］−１３．１×［Ｖ］−１０．３×［Ｍｏ］ ・・・（２） Al: 3.0-5.0 mass%,
V: 1.0-3.0 mass%,
Fe: 1.0-1.8 mass%,
Mo: 0.9 to 1.7 mass%,
O: 0.05-0.25 mass% and
N: 0.02 to 0.15 mass% ,
The balance consists of Ti and inevitable impurities, the aluminum equivalent Al eq represented by the following formula (1) is 4 to 8 mass%, and the β transformation temperature T β represented by the following formula (2) is 880 to 980 ° C. α-β type titanium alloy.
Al eq (mass%) = [Al] + 10 × [O] + 27.7 × [N] (1)
T β (° C.) = 886 + 147.7 × [O] + 294.3 × [N] + 20.4 × [Al] −19.8 × [Fe] −13.1 × [V] −10.3 × [Mo ] (2)
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| CN113967711A (en) * | 2021-10-22 | 2022-01-25 | 西安赛特思迈钛业有限公司 | Hot processing method of two-phase titanium alloy bar |
| CN114807678B (en) * | 2022-04-28 | 2023-09-08 | 中国科学院金属研究所 | High-strength high-toughness weldable high-temperature titanium alloy and preparation method thereof |
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| JPH0823053B2 (en) * | 1989-07-10 | 1996-03-06 | 日本鋼管株式会社 | High-strength titanium alloy with excellent workability, method for producing the alloy material, and superplastic forming method |
| JP2541042B2 (en) * | 1991-08-30 | 1996-10-09 | 日本鋼管株式会社 | Heat treatment method for (α + β) type titanium alloy |
| JP2737498B2 (en) * | 1991-12-24 | 1998-04-08 | 日本鋼管株式会社 | Titanium alloy for high density powder sintering |
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| CN112795799A (en) * | 2020-11-23 | 2021-05-14 | 昆明理工大学 | Cold cathode EB furnace smelting method of forging-free direct rolling Ti-Al-V-Fe alloy ingot |
| CN112795799B (en) * | 2020-11-23 | 2021-11-26 | 昆明理工大学 | Cold cathode EB furnace smelting method of forging-free direct rolling Ti-Al-V-Fe alloy ingot |
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